Nickel-based superalloys are known. At high temperatures, single-crystal components made of these alloys have inter alia a very good material strength but also good corrosion and oxidation resistance, as well as a good creep strength. Owing to these properties, when using such materials for example in gas turbines, the intake temperature of the gas turbines can be raised so that the efficiency of the gas turbine system increases.
Simply speaking, there are two types of single-crystal nickel-based superalloys.
The first type, to which the present invention also relates, can be fully heat treated (solution annealed) so that the entire γ′ phase lies in solution. This is for example the case for the known alloy CMSX4 with the following chemical composition (data in wt. %): 5.6 Al, 9.0 Co, 6.5 Cr, 0.1 Hf, 0.6 Mo, 3 Re, 6.5 Ta, 1.0 Ti, 6.0 W, remainder Ni or the alloy PWA 1484 with the following chemical composition (data in wt. %): 5 Cr, 10 Co, 6 W, 2 Mo, 3 Re, 8.7 Ta, 5.6 Al, 0.1 Hf and the known alloy MC2 which, in contrast to the previously mentioned alloys, is not alloyed with rhenium and has the following chemical composition (data in wt. %): 5 Co, 8 Cr, 2 Mo, 8 W, 5 Al, 1.5 Ti, 6 Ta, remainder Ni.
A typical standard heat treatment for CMSX4 is for example as follows: solution annealing at 1320° C./2 h/shield gas, rapid cooling with a blower.
The second type of single-crystal nickel-based alloys is not fully heat treatable, i.e. in this case the entire proportion of the γ′ phase does not enter solution during solution annealing, rather only a certain part of it does. This is for example the case for the known superalloy CMSX186 with the following chemical composition (data in wt. %): 0.07 C, 6 Cr, 9 Co, 0.5 Mo, 8 W, 3 Ta, 3 Re, 5.7 Al, 0.7 Ti, 1.4 Hf, 0.015 B, 0.005 Zr, remainder Ni and the alloy CMSX486 with the following chemical composition (data in wt. %): 0.07 C, 0.015 B, 5.7 Al, 9.3 Co, 5 Cr, 1.2 Hf, 0.7 Mo, 3 Re, 4.5 Ta, 0.7 Ti, 8.6 W, 0.005 Zr, remainder Ni.
The nickel-based superalloys of the second type are usually exposed to a two-stage heat treatment (ageing process at lower temperatures) since at high temperatures, such as are typically used for solution annealing the alloys of the first type, the melting point initiation temperature is already reached and the alloy therefore undesirably begins to melt.
A typical two-stage heat treatment of the alloy CMSX186 is for example as follows:                1st stage: 1080° C./4 h/blower        2nd stage: 870° C./20 h/blower        
The creep strength of the first type of nickel-based superalloys is normally higher than that of the second type, assuming that the alloys belong to the same generation. This is primarily due to the fact that the dissolved γ′ is the main source of the achievable strength.
Nickel-based superalloys for single-crystal components, such as are known from U.S. Pat. No. 4,643,782, EP 0 208 645 and U.S. Pat. No. 5,270,123, contain mixed-crystal strengthening alloy elements, for example Re, W, Mo, Co, Cr, and γ′ phase-forming elements, for example Al, Ta and Ti. The content of high-melting alloy elements (W, Mo, Re) in the basic matrix (austenitic γ phase) increases continuously with the rise in the working temperature. Thus, for example, conventional nickel-based superalloys for single crystals contain 6-8% W, up to 6% Re and up to 2% Mo (data in wt. %). The alloys disclosed in the documents cited above have a high creep strength, and good LCF (Low Cycle Fatigue) and HCF (High Cycle Fatigue) properties as well as a high oxidation resistance.
These known alloys were developed for aircraft turbines and therefore optimized for short- and medium-term use, i.e. the working time is configured for up to 20,000 hours. In contrast to this, industrial gas turbine components must be configured for a working time of up to 75,000 hours or even more.
After a working time of 300 hours during experimental use in a gas turbine at a temperature above 1000° C., for example, the alloy CMSX-4 of U.S. Pat. No. 4,643,782 shows a large growth of the γ′ phase which is detrimentally associated with an increase in the creep strength of the alloy.
Another problem with the known nickel-based superalloys, for example the alloys known from U.S. Pat. No. 5,435,861, is that the castability for large components, for example in gas turbine blades with a length of more than 80 mm, leaves much to be desired. It is extremely difficult to cast a perfect, relatively large directionally solidified single-crystal component from a nickel-based superalloy because most of these components comprise defects, for example small-angle grain boundaries, “freckles” i.e. defects due to a series of co-oriented grains with a high eutectic content, axial variations, microporosities etc. These defects weaken the components at high temperatures, so that the desired lifetime or operating temperature of the turbine is not achieved. But since a perfectly cast single-crystal component is extremely expensive, the industry tends to allow as many defects as possible without compromising the lifetime or operating temperature.
Some of the most frequent defects are grain boundaries, which are particularly detrimental to the high-temperature properties of the single-crystal articles. While small-angle grain boundaries have only a comparatively minor effect on the properties for small components, they are of great relevance for the castability and the oxidation behavior at high temperatures in the case of large SX or DS components.
Grain boundaries are regions of high local disorder in the crystal lattice since the neighboring grains meet in these regions and there is therefore a certain disorientation between the crystal lattices. The greater this disorientation is, the greater is the disorder i.e. the greater is the number of dislocations in the grain boundaries which are needed so that the two grains fit together. This disorder is in direct correlation with the behavior of the material at high temperatures. It weakens the material when the temperature is increased above the equicohesive temperature (=0.5×melting point in K).
This effect is known from GB 2 234 521 A. At a testing temperature of 871° C. for a conventional nickel-based single-crystal alloy, for example, the breaking strength decreases greatly when the disorientation of the grains is more than 6°. This has also been found for single-crystal components with a directionally solidified structure, so the opinion has generally been accepted that disorientations of more than 6° are not permissible.
The tolerance with respect to a deviation in the small-angle grain boundaries, or grain boundary disorientation, is generally greater for the second type of nickel-based superalloys than for those which are not fully heat treatable.
It is also known from the cited GB 2 234 521 A that enriching nickel-based superalloys with boron or carbon in the case of directional solidification generates structures which have an equiaxial or prismatic grain structure. Carbon and boron strengthen the grain boundaries since C and B cause the precipitation of carbides and borides, which are stable at high temperatures, on the grain boundaries. The presence of these elements in the grain boundaries and along the grain boundaries furthermore reduces the diffusion process, which is a main cause of the grain boundary weakness. It is therefore possible to increase the disorientations to as much as 10° to 12°, and nevertheless achieve good properties of the material at high temperatures.
Particularly for large single-crystal components made of nickel-based superalloys, however, these small-angle grain boundaries detrimentally affect the properties. Furthermore, microalloying with B and C is limited to a few hundred ppm C and only a few ppm B, since said elements on the one hand have only a low solubility in the matrix and on the other hand have a strong effect on the undesired reduction of the initial melting point of this alloy.
Heat treatment methods for nickel-based superalloys are known from US 2004/0055669 A1, EP 0 155 827 A2, WO 2004/038056 A1 and DE 196 17 093 A1, in which the alloy is only partially solution annealed in a first heat treatment step and a two-stage ageing treatment known per se is carried out at respectively lower temperatures in a second step.